Silicon nitride having a high tensile strength

ABSTRACT

A ceramic body comprising at least about 80 w/o silicon nitride and having a mean tensile strength of at least about 800 MPa.

STATEMENT OF GOVERNMENT SUPPORT

This invention was developed under U.S. Government Contract No.DE-AC05-84OR21400 awarded by The United States Department of Energy.

This is a divisional of application Ser. No. 08/324,937 filed on Oct.18, 1994, now U.S. Pat. No. 5,571,760.

BACKGROUND OF THE INVENTION

Advanced structural ceramic materials have gained the attention ofindustry by virtue of their superior performance qualities. Thesequalities, such as superior high temperature strength, high toughness,resistance to thermal shock and resistance to oxidation provide thebasis for their potential use in a variety of applications.

Despite their enormous potential, advanced structural ceramics ingeneral and silicon nitride in particular have yet to capture manymarkets, principally due to the perception that they tend towardcatastrophic failure, and thus are unsuitable for uses in applicationsrequiring high reliability.

While significant progress has been made in the development of strong,tough, refractory ceramics for demanding applications such as the AGTrotor, processing these materials sometimes still introducesstrength-limiting and reliability-reducing flaws.

One measure of the extent of the presence of flaws is the tensilestrength of the ceramic. Typically, a conventional ceramic materialoften possesses a tensile strength in the neighborhood of only betweenabout 50% and 90% of its flexural strength. Accordingly, the averagetensile strength of conventional silicon nitride materials is often inthe range of about 650 to 750 MPa.

Accordingly, it is the object of the present invention to provide asilicon nitride ceramic having an average tensile strength of at leastabout 800 MPa.

SUMMARY OF THE INVENTION

In accordance with the present invention, there is provided a siliconnitride ceramic having an average tensile strength of at least about 800MPa, preferably about 900 MPa, more preferably about 1000 MPa. Inpreferred embodiments, there is provided a process for producing a hightensile strength silicon nitride ceramic, comprising the steps of:

a) milling a batch of silicon nitride powder in water to produce aslurry having a particle size distribution characterized by a d10 ofabout 0.28 microns, a d50 of about 0.67 microns, and a d90 of about 1.27microns,

b) passing the slurry through a filter having a pore size of betweenabout 2 and about 10 microns,

c) concentrating the slurry to a solids content of between about 70 andabout 72 w/o,

d) passing the slurry through a magnetic separator,

e) casting the slurry at a casting rate slope of less than about 10 mm²/min to produce a green body comprising at least about 80 w/o siliconnitride,

f) drying the green body to less than about 3 w/o water,

g) degassing the green body,

h) sintering the green body to yield a dense ceramic comprising at leastabout 80 w/o silicon nitride,

i) machining the dense ceramic, and

j) heat treating the dense ceramic to yield a high tensile strengthsilicon nitride ceramic.

BRIEF DESCRIPTION OF THE FIGURES

FIG. 1 presents a plot of cast thickness over time produced from aprogrammed pressure profile of the present invention.

FIG. 2 presents a diagram of the milling operation of the presentinvention.

FIG. 3 presents the tensile strength data of the present invention alongwith 2-parameter and 3-parameter Weibull correlations thereof.

FIG. 4 presents a toughness-aspect frequency correlation characterizingthe ceramics of the present invention.

DETAILED DESCRIPTION OF THE INVENTION

Without wishing to be tied to a theory, it is believed that theprocessing of the present invention results in a silicon nitride bodyhaving a very low frequency of flaws. Flaws can be characterized asintrinsic (i.e., flaws occurring randomly within a conventionalprocessing cycle) or extrinsic (i.e., flaws unrelated to conventionalprocessing). Surface glass blisters and volume microstructural flaws(i.e., inclusions, agglomerates and porosity) are considered to beintrinsic flaws since blisters result from oxidation treatments andmicrostructural flaws arise from imperfect mixing and contamination.Machining damage is considered to be an extrinsic defect since it arisesfrom poorly controlled machining and is not considered part ofconventional powder processing, forming or densification.

Again, without wishing to be tied to a theory, it is further believedthat the process of the present invention provides a controlled methodof producing reliable silicon nitride ceramics by reducing the frequencywith which the above-cited flaws arise. In particular, blistergeneration can be minimized by controlling the oxidizing environment.Porosity can be controlled by deairing the slurry. Inclusions can belowered by utilizing a process including: a) a Class 10,000 clean room,b) separation of dry powder from the milling area, c) filtered anddeionized water, d) magnetic separation, e) filtration of the milledslurry, f) a closed system from milling to casting, g) non-metalliccontact surfaces for equipment contacting the silicon nitride slurry,and h) high purity/low attrition hipped milling media. Agglomerates canbe reduced by using a fresh, sonicated slip. Finally, machining damagecan be controlled by improved machining procedure.

The silicon nitride starting materials may be any silicon nitride powderor precursor thereof typically used in the processing of fine ceramics.However, since it is the object of the present invention to produce aceramic relatively free of flaws, it is desirable to use highly puresilicon nitride starting materials. Highly pure starting materialsinclude chemically derived silicon nitride powders and preferably,purified native silicon nitride powders. In silicon nitride startingmaterials of high purity, iron and residual chlorine contamination inthe powder should be no more than about 30 ppm and 50 ppm, respectively.There should be no detectable free silicon. More preferably, high puritysilicon nitride powder available from Ube Industries of Tokyo, Japan isselected as the silicon nitride starting material.

The oxygen content of the silicon nitride starting materials istypically between about 2.5 w/o and about 3.0 w/o of the silicon nitridestarting material. If the oxygen content is no more than about 2.5 w/o,it is typically inhomogeneously distributed within the silicon nitridestarting material. Accordingly, these starting materials are oftensubject to thermal oxidations to produce a more homogeneous oxygendistribution. These treatments typically raise the oxygen content toabout 3 w/o.

In especially preferred embodiments, the silicon nitride startingmaterial is a powder blend comprising a first portion having an averageparticle size of between about 0.60 and about 2.3 microns, a secondportion having an average particle size of between about 0.30 and about1.60 microns, and a third portion having an average particle size ofbetween about 0.20 and about 1.20 microns. This particular blend ispreferred because it allows packing density to be easily controlled andoptimized. Typically, the first portion is Ube E03, the second portionis Ube E05, and the third portion is Ube E10. Powder blends disclosed inU.S. Pat. No. 5,001,091 may also be suitably used as the startingmaterials of the present invention.

The phase of the silicon nitride starting material may be alpha, beta orany combination thereof. Preferably, between about 92 w/o and about 98w/o of the silicon nitride starting material is in the alpha phase. Inmore preferred embodiments, it is about 95 w/o of the silicon nitridestarting material.

The isoelectric point of the silicon nitride starting powder (IEP) is animportant process control parameter for the present invention because itcontrols the casting rate and hence the density and uniformity of thecast body. The isoelectric point of the starting powder should be withinthe range of about 6.8 to about 7.2 pH units. When the IEP of the powderis outside this range, undesirable casting rates may be realized,resulting in undue density gradients in the cast body.

The production of highly reliable ceramics requires highly homogeneousstarting materials. Typically, however, starting materials showsignificant lot-to-lot particle size variability, with the surface areaof starting material powders often varying by about 10% to about 40%.Since there is often a significant variability in starting materialparticle size, the raw powder lots are milled to achieve a consistentparticle size distribution prior to green forming. Depending upon thecondition of the as-received powder, milling conditions may have to bevaried to arrive at the desired size distribution. Any millingtechnique, including ball milling, attrition milling and vibrationmilling, may be used to achieve this purpose. However, since vibrationmilling provides for continuous monitoring and sampling throughout themilling operation and its low impact forces allow for the milling of lowsolids content mixes without excessive media wear, it is the preferredmode of milling. In some embodiments, milling is accomplished in a SwecoM18 vibratory mill, manufactured by Sweco, Inc. of Florence, Ky.Preferably, this mill is charged with about 43338 cc of deionized waterand about 76 kg of high purity hipped silicon nitride milling mediahaving 1/2 inch cylindrical bodies with hemispherical ends. The sizedistribution of milled powders of the present invention is typicallycharacterized by 3 parameters: d10, d50 and d90. For the purposes of thepresent invention, in a powder having a dA of X microns, A % of theparticles have a diameter of no more than X microns. The milled powdersof the present invention typically have a d10 of between about 0.24 andabout 0.32 microns, preferably about 0.28 microns; a d50 of betweenabout 0.64 and about 0.70 microns, preferably about 0.67 microns; and ad90 of between about 1.17 and 1.37 microns, preferably about 1.27microns.

In some embodiments, the silicon nitride milling media becomes acomponent of the milled product due to attrition during milling. Thiscomponent typically represents between about 0.2 and about 1.0 w/o ofthe final milled product. The effect of this attrition upon the productof the present invention is typically some nominal increase in cationimpurities (i.e. an increase of less than about 100 ppm in cationimpurities).

Although a non-aqueous liquid milling medium such as an alcohol may beused in the milling step of the present invention, water is preferablyused. It has been found that water milling not only results in about a30% reduction in milling time as compared to conventional alcoholmilling, but also results in less inclusion contamination since there isno need to dry and redisperse (and thereby expose to contamination) themilled powder in water prior to casting. When water is selected as theliquid milling media, however, the potential for oxidation of thesilicon nitride starting materials must be considered. Accordingly, itis desirable to use a surfactant in order to enhance powder dispersionand to provide a protective coating on the powder surface whichminimizes oxidation of the powder.

Because silicon nitride does not easily sinter by itself, it may bedesirable to add sintering aids to the silicon nitride startingmaterials of the present invention. Rare earth oxides such as yttria maybe used in amounts typically used in the art to assist sintering. Insome embodiments, yttria powder, manufactured by Molycorp of WhitePlains, N.Y., is added during the milling operation, preferably in anamount of about 4 w/o of the silicon nitride powder. During co-milling,some amount of undesirable foaming may be observed. Accordingly, about0.0025 w/o of a defoaming agent such as Mazu DF204, available from PPGIndustries of Gurnee, Ill., can be used to mitigate the foaming.

Because the above-described milling operation may not only produceagglomerated silicon nitride starting materials but also introducecontamination into the milled product, it is often desirable to pass themilled product through filters to eliminate these newly-generatedagglomerates and contaminants. Preferably, these filters have pore sizesof between about 2 and about 5 microns. Although any filterconventionally used in the filtration of silicon nitride slurries may beused, a cartridge polymer filter is preferred. More preferably, thefilter is a 5 micron RIF050 Pall Profile Filter, available from ChisolmCorp. of Cranston, R.I. Typically, filtration can easily removeparticles greater than about 10 microns when the milled powder primarilycomprises particles less than about one micron. However, when filtrationof smaller particles (i.e. less than about 10 microns) is desired, ithas been found that the milled product should be diluted with about 30w/o of deionized water prior to filtration.

Because iron inclusions often reduce the reliability of silicon nitrideceramics, the process of the present invention typically includes amagnetic separation step. In preferred embodiments, the filtered productin the form of a slurry is recirculated through a magnetic separator. Inmore preferred embodiments, the magnetic separator is Model 43XPFerrofilter, manufactured by S. G. Frantz of Trenton, N.J. The slurry isusually fed through the magnetic separator at a rate of 2 gallons perminute and recirculated for one hour. There should be no metallicinclusions found in powder dried immediately after it is run through themagnetic separator, as determined by a microfocus x-ray having adetection limit of about 25 microns.

After the slurry is subjected to magnetic separation, it usually has asolids loading of about 30 w/o. Because such a slurry is likely to betoo dilute to be effectively cast, a concentration step is typicallyundertaken. In preferred embodiments of the present invention, theslurry is concentrated to a solids loading of between about 65 and about74 w/o. Above about 74 w/o, the slurry is thought to be susceptible todilatency. In preferred embodiments, the concentration step is carriedout with a crossflow filter, as described in U.S. Pat. No. 5,229,339,the specification of which is wholly incorporated by reference. Thecross flow filter allows the slurry to be concentrated to high (i.e.,65-74 w/o) levels while maintaining suspension rheology (30-70 cP)suitable for casting.

Slips concentrated in accordance with the present invention maysometimes contain agglomerates which can act as flaws. One method ofbreaking up these agglomerates involves sonicating the slip. Thus, inpreferred embodiments, the concentrated slip is sonicated with a 5 kWultrasonic horn for five minutes. The sonication produces adeagglomerated slurry.

Slips processed in accordance with the present invention may alsocontain air pockets. These air pockets remain throughout casting andproduce strength limiting porosity when the cast is sintered.Accordingly, in some embodiments, the sonicated slip is deaired byplacing it in a 60 mm Hg vacuum for about 5 minutes.

Once deaired, the slip typically has a viscosity of about 70 cp and a pHof about 9.5. Preferably , the slip has a viscosity of about 60 cp, a pHof about 9.0, and an isoelectric point (IEP) of about 6.8. However, aslip is considered suitable if it has a viscosity of about 60-70 cp anda pH of about 9-10. Because all of these parameters change as the slipages, care must be taken to process the slip before the parameterssignificantly change. If excessively aged slip is used, agglomeratestypically form. It has been found that the slip should be used withinabout 5 days of its deairing, preferably within about 3-5 days, in orderto prevent agglomerate formation.

In casting green pieces, it is important to have a relatively slow cast.Excessive casting rates provide inadequate time for the optimum packingof the powder particulates during filtration, leading to both low andinhomogeneous cast densities. Accordingly, the slurry of the presentinvention may be further characterized by its casting rate slope("CRS"). The CRS of the slurry is defined as the slope of the square ofthe thickness of a cast body versus time when the slurry is cast under aset of standard conditions. In some embodiments of the presentinvention, the standard conditions were set at about 200 g of slurrycast at a pressure of about 0.21 MPa onto a Plaster of Paris porousmaterial to produce a 3 inch diameter disc. When the CRS is less thanabout 10 mm² /min, the cast rate is suitably slow to provide a cast bodyof high and homogeneous density. Typically, the slurry concentrationrequired for casting is between about 70 w/o and about 72 w/o solids.

In some embodiments of the present invention, slower casting rates(i.e., casting rates between about 2 and about 5 mm² /min) have yieldedhighly reliable ceramic bodies which do not show excessive warping evenwhen the ceramic has a high aspect ratio. It is believed the slowercasting rates may be achieved by using a slurry having a solids contentof between about 57 and about 63 w/o, preferably about 60 w/o, or aslurry whose d10 is between about 0.25 and about 0.30 microns. Thus, inaccordance with the present invention, there is provided an unmachinedceramic body comprising at least about 80 w/o silicon nitride, saidceramic body having a mean tensile strength of at least about 800 MPa, arunout of less than about 30 mil per about 6.8 inches, wherein runout isdefined as the depth of bend from a perfectly straight configuration,and an aspect ratio of at least about 3:1.

Once the slip properties (including its CRS) are found to be suitable,the slip is ready to be cast. In preferred embodiments of the presentinvention, the casting mold is a two piece Plaster of Paris mold havingan average pore size of about 2 microns which has been soaked indistilled water for about 2 minutes. In some embodiments, one piece ofthe mold is Plaster of Paris and the other piece of the mold is plastic.Preferably, the mold has a single casting front. It was found that moldsproviding a filtration surface wholly around the casting cavity (andthus providing multiple casting fronts) often produced cast pieces whichcracked upon drying at the points where casting fronts met. Conversely,molds designs which present a single casting front yielded cast barswhich showed no cracking even under uncontrolled, i.e. ambient, dryingconditions. Tomographic examination of bars cast under multiple frontsshowed a low density at the center of the bars. Conversely, the centersof those bars cast with a single front showed high density.

Typically, green bodies are cast under a constant applied pressure ofbetween about 0.1 and about 0.5 MPa. However, permeability changes oftenarise which threaten the uniform density of a cast body. Accordingly, avariable pressure profile is preferably used to accommodate thepermeability change. See FIG. 1. The use of such a profile results in alinear casting rate and hence a uniform density throughout the castbody.

Once the slurry is cast in accordance with the present invention, theresulting green body may be optionally dried. Generally, any standardmethod of drying may be used, including controlled humidity drying. Thehumidity of the drying step may be controlled or uncontrolled.Preferably, the dried casting contains no more than about 3 w/o water.

The dried casting may be subjected to a degassing step to remove allchemisorbed species. When degassing is desired, the green body istypically placed in a furnace having a temperature of about 1450° C. anda nitrogen atmosphere for about 60 minutes.

After the casting is degassed, it is sintered to a dense ceramic. Anymethod of sintering known in the art can be utilized, including hotpressing, pressureless sintering, gas pressure sintering, and hotisostatic pressing ("hipping"). Moreover, conventional sintering timesand temperatures may be used. In preferred embodiments of the presentinvention, the casting is subjected to glass encapsulated hipping inaccordance with U.S. Pat. No. 4,446,100. More preferably, the sinteringcycle comprises an initial soak of 1840° C. for 20 minutes and a secondsoak of 1740° C. for about 45 minutes, each soak undertaken at 30,000psi.

After the sintered ceramic is produced, it often contains a reactionlayer and slight dimensional distortion. Hence, machining of the ceramicis usually necessary. In preferred embodiments of the present invention,machining is carried out as per ASTM STANDARD C 1161-90. In other words,the preferred machining process involves wet grinding between centersusing a computer numerical control ("CNC") grinding machine, a roughgrinding step and a finish grinding step. In the rough grinding step, aresin bonded 180 grit diamond wheel is used to grind the specimen at a0.001" per pass feed rate, leaving 0.005" stock for finish grinding. Inthe finish grinding step, a resin bonded, 320 grit diamond wheel is usedto grind the specimen at a 0.0002" per pass feed rate. In each step, thewheel speed is 5500-7500 surface feet per minute ("SFM") and the workspeed is 50-200 rpm. The resulting ceramic measures a surface roughnessof less than about 8 microinches.

Machining of the ceramics of the present invention often producessubsurface flaws in the ceramic typically appearing as half-penny shapedcracks. Accordingly, the machined ceramic must often be heat treated toanneal these flaws. Heat treating the machined ceramic typicallyrequires subjecting it to a temperature of about 1000° C. in an airatmosphere for about 20 hours.

When the process controls described above are incorporated in siliconnitride processing, the resulting ceramics typically have a mean tensilestrength of at least about 800 MPa, preferably at least about 900 MPa,and more preferably at least about 1000 MPa, with an associated WeibullModulus of between about 10 and about 28, a mean flexural strength ofbetween about 900 and about 1000 MPa, and a toughness of between about5.5 and about 6.4 MPa m^(1/2). In some embodiments, the tensile strengthof a ceramic of the present invention is characterized by a 3-parameterWeibull Modulus having a cutoff of at least about 665 MPa, and istypically at least about 95% to 105% of its flexural strength.

The microstructure of these ceramics typically contain less than about0.4 v/o porosity, less than about 8 v/o inclusions, and less than about17 v/o agglomerates. Further, the inclusions are typically less thanabout 25 microns in length while the agglomerates are typically lessthan about 20 microns in diameter.

Between about 20 v/o and about 30 v/O of the silicon nitride grains aretypically alpha phase, the remainder being in the beta phase. It wasobserved that when between about 20 w/o and about 30 w/o of the siliconnitride grains in the ceramic are in the alpha phase, toughness isenhanced. Without wishing to be tied to a theory, it is believed that anabbreviated sintering cycle produces this phenomenon. In typicalsintering cycles, alpha phase grains transform to long, thin, highaspect ratio beta phase grains which then coarsen to lower aspect ratiosas sintering progresses. Moreover, transformation does not occur in eachgrain simultaneously. Thus, transformed grains typically coarsen beforeall alpha phase grains are completely transformed. Since conventionalsintering theory teaches to continue sintering until all the alpha phasegrains are transformed, typical sintering cycles produce ceramic havingpredominantly low aspect ratio beta grains and essentially no alphagrains. However, it has been found that high aspect ratio beta grainsare superior toughening agents to lower aspect ratio coarse beta grains.Accordingly, it is believed that abbreviated sintering schedules whichproceed far enough to transform alpha grains to high aspect ratio betagrains, but not so far as to substantially coarsen the beta grains, willprovide an enhanced toughness material. Such abbreviated sinteringcycles typically include a soak stage wherein the green body is hippedat a temperature of between about 1800° and about 1840° C. for betweenabout 15 and about 30 minutes, preferably at a pressure of about 30,000psi. A second soak stage may optionally be added wherein the green bodyis hipped at a temperature of between about 1700° and about 1780° C. forbetween about 30 and about 180 minutes, preferably at between about1740° and about 1760° C. for between about 30 and about 60 minutes. Ithas been observed that such abbreviated cycles which maximize thefrequency of high aspect ratio beta grains typically fail to transformbetween about 20 and about 30 w/o of the alpha grains present in thepresintered body. In the ceramics produced from such abbreviated cycles,at least about 80% the beta grains have a thickness less than about 1micron and at least about 95% of the beta grains typically have anaspect ratio of at least about 4.

EXAMPLE I

About 43338 g of deionized water containing about 1 w/o AG surfactant,manufactured by Cavcomod of Woonsocket, R.I., was added to asemi-continuous polypropylene holding tank and vortex-mixed for 5minutes to produce a well-stirred dispersion. To this dispersion wasadded about 1156 g of 5603X High Surface Area yttria, manufactured byMolycorp of White Plains, N.Y.; about 13868 g of Ube E03 siliconnitride; about 8320 g of Ube E05 silicon nitride; and about 5548 g ofUbe E10 silicon nitride. The resultant slurry was then mixed in the tankfor about 1 hour at high speed. After the 1 hour mixing cycle, theslurry was transferred to a Sweco M18 chamber, manufactured by Sweco,Inc. of Florence, Ky., containing about 76,000 g of half-inch hippedsilicon nitride media. See FIG. 2. The transfer was undertaken byconstructing a flow circuit from the tank through a pump to the Swecomill and back to the tank. The tank, pump and Sweco Mill were connectedby plastic tubing, an inert plastic tubing having a half-inch innerdiameter. Each unit operation shown in the flow circuit of FIG. 2 wasconnected by this tubing. Sufficient slurry was initially pumped fromthe tank to fill the Sweco mill to about the middle of the fill port.When the mill was so filled, the pumping ceased, the stainless steelcover of the mill was secured, the mill was turned on, and the pump wasrestarted. The slurry was milled therein for about 1 hour.

At this point, the slurry was well-dispersed and deagglomerated. Theflow circuit was reconstructed so that the slurry could flow from thetank through a pump to a Model 43XP ferro-filter, manufactured by S. G.Frantz Co. of Trenton, N.J., to the mill to the tank. About 29 g ofammonium citrate (adjusted to a pH of 9.5 with a solution of 70%ammonium hydroxide and 30% deionized water) and about 0.75 g of MazuDF204 were then added to the slurry. The pH of the tank was thenadjusted to about 9.6, using ammonium hydroxide. The slurry wasrecirculated through the circuit, thus passing through the ferrofilter,for about one hour. The flow was then redirected to its original pattern(i.e., tank to pump to mill to pump) and milling continued.

Samples of the slurry were periodically taken from the circuit in orderto determine the mean particle size. When the mean particle size wasabout 0.67 microns, the milling ceased, the pump was turned off, thestainless steel mill lid was removed, and the slurry was drained intothe holding tank.

After the slurry was transferred to the holding tank, the flow circuitwas broken between the holding tank and the ferro filter. To the ferrofilter exit side was attached, sequentially, a Pall filter cartridgehaving a 5 micron filter and a Ceraflo cross-flow tank, manufactured byMillipore Corporation of Bedford, Mass. After the cartridge was wettedwith deionized water, the pump was turned on, and flow proceeded fromthe tank to the pump to the ferrofilter to the Pall filter to thecrossflow tank.

Next, the slurry was transferred into the cross flow tank, and its pHwas adjusted to about 9.7 by ammonium hydroxide addition. The diaphragmpump regulator of the crossflow tank was set at about 80 psi and thepump pulsed at about 1 cycle per second. The top filtrate outlet of thefilter tube was opened and clear filtrate exited from the outlet at aninitial rate of at least about 1 liter per minute. Filtrate flowdecreased as concentration proceeded. After the entire slurry was pumpedinto the crossflow tank, pump speed increased to about 2 cycles persecond and maintained there for about 3 hours.

When the slurry concentrated to about 60 w/o solids and had a specificgravity of about 1.7 g/cc, casting rate tests were performed todetermine the CRS of the slurry. Based upon the CRS value obtained, atarget slurry concentration was calculated which would result in a CRSof less than about 4 mm² /min. Typically, the target slurryconcentration was calculated to be between about 70 w/o and about 72 w/osolids. These target concentrations were achieved by further crossflowfiltration.

When the desired slurry concentration was obtained, the crossflow pumpwas shut off. The hose from the tank outlet was disconnected and theconcentrated slurry was pumped into the pressure cast system holdingtank. The slurry was kept circulating in the pressure cast systemholding tank by using a diaphragm pump (capacity of about 13 kg/m and amaximum pressure of about 100 psig) at a speed of about 2 cycles persecond. The pump speed was decreased so that a vortex could not form inthe center of the casting tank as the volume of the slip decreasedduring casting.

The concentrated slurry was then cast in a two piece, single frontPlaster of Paris mold having an average pore size of about 2 micronswhich was soaked in distilled water for about two minutes.

Next, the green body was placed in a furnace having a temperature ofabout 1450° C. and an atmosphere of nitrogen for about 60 minutes todegas it.

After the green body was degassed, it was subjected to glassencapsulated hipping in accordance with U.S. Pat. No. 4,446,100 havingan initial soak of 1840° C. for 20 minutes and a second soak of 1740° C.for about 45 minutes, each soak at 30,000 psi.

After the sintered ceramic is produced, machining is carried out usingspecifications from ASTM STANDARD C 1161. The machining process involvedwet grinding between centers using a CNC grinding machine and twomachining steps. In the first step, a resin bonded 180 grit diamondwheel was used to grind the specimen at a 0.001" per pass feed rate,leaving 0.005" stock for finish grinding. In the second step, a resinbonded, 320 grit diamond wheel was used to grind the specimen at a0.0002" per pass feed rate. In each step, the wheel speed was 5500-7500SFM and the work speed was 50-200 rpm. The resulting ceramic measured asurface finish of less than about 8 microinches.

After the machining was completed, the machined ceramic was subjected toa temperature of about 1000° C. in an air atmosphere for about 20 hours.

Approximately 320 tensile bars were made in accordance with thisExample. The final machined cylindrical buttonhead tensile bars haddimensions set forth in Gerber et al., "Rotor Data Base Generation"CTAHE Semiannual Progress Report for October 1990 through March 1991,pp. 340-361, July 1991, except that the gage diameter was 6.0+/-0.1 mm.The mechanical properties and microstructures of these ceramics wereanalyzed.

Toughness was measured in accordance with indentation techniquesdescribed in Chantikul et al. "A Critical Evaluation of IndentationTechniques for Measuring Fracture Toughness II: Strength Methods" J. Am.Cer. Soc. 64(9), pp. 539-544 (1981). Toughness was found to be in therange of between about 6.0 to 6.4 MPa m^(1/2). The average toughness wasabout 6.2 MPa m^(1/2).

Tensile strength was measured as follows: The cylindrical buttonheadtensile specimens were tested at room temperature on an Instrom Model8562, manufactured by Instrom Corp. of Canton, Mass., utilizing theInstrom "Super-grip" and straight tri-split copper collet system. Adouble ramp loading procedure was used to test all specimens. Thespecimen was initially loaded to about 6668N (1500 lbs) at about 39MPa/min (250 lbs/min). This allowed time for the fully annealed coppercollets to deform to the buttonhead radius of the specimen. After theinitial ramp to about 6668N, the specimen was then loaded to failure ata stressing rate of about 600 MPa/min. The load train was checked beforetesting with an alignment tool for actuator/load cell alignment and astrain gaged tensile specimen for Instrom "Supergrip" alignment.

For the 320 tensile bars which fractured within about one cubiccentimeter uniformly stressed gage section, the mean tensile strengthwas about 997 MPa. The range of the data was from a low value of 540 MPato a peak value of 1237 MPa. More than 80% of the 320 tensile barsdescribed above had a tensile strength of at least about 900 MPa. About50% had tensile strength values in excess of 1 GPa.

SEM fractography analysis for this sample set of 320 tensile barsdetermined the strength distribution to be multimodal with several flawpopulations contributing to the strength-defining fast fractures.

A competing risk analysis was undertaken to determine whether theintrinsic tensile strength data could be better characterized by a2-parameter or 3-parameter Weibull fit. See FIG. 3. The confidence level(R²) for the 2-parameter model was determined to be 95.79% while that ofthe 3-parameter model was 99.29%. Thus, the competing risk analysissuggested that the strength distribution associated with processrelated, intrinsic strength impairing flaws is best represented by a3-parameter Weibull model. Approximately one half of the data fall intothis "intrinsic" category while the other half involved machining damagerelated failure origins. Set forth below in TABLE 1 are the intrinsicstrength data for the 3-parameter Weibull model which exclude failuresdue to machining damage. As shown in Table 1, the 3-parameter modelyields an intrinsic strength threshold equal to about 665 MPa.

                  TABLE 1                                                         ______________________________________                                                       2-Parameter                                                                             3-Parameter                                                         Weibull Model                                                                           Weibull Model                                        ______________________________________                                        Mean Tensile Strength (MPa)                                                                    1058        1067                                             Characteristic Tensile                                                                         1102        1109                                             Strength (MPa)                                                                Shape/Modulus    13          4                                                Threshold (MPa)  0           665                                              R.sup.2          95.79%      99.29%                                           ______________________________________                                    

The microstructure of the ceramics made in accordance with Example 1were also examined. XRD analysis indicated that between about 20 w/o andabout 30 w/o of the silicon nitride grains were alpha, the remainderbeing beta. Without wishing to be tied to a theory, it is believed thatthe relatively low sintering temperature combined with a relativelyshort soak time prevented a higher degree of conversion of the alphasilicon nitride to the beta phase.

Next, grain thicknesses of the beta silicon nitride grain were estimatedby the linear intercept method. It was found that at least about 80% ofthese grains had a thickness of less than about 1 micron.

Third, the average aspect ratio of the silicon nitride grains wereinvestigated. The length and thicknesses of beta grains were estimatedby the linear intercept method. The average aspect ratio was about 6.The high aspect ratio grains (i.e., grains having an aspect ratiogreater than eight) tended to have thicknesses less than about 0.6microns.

Fourth, the frequency of high aspect ratio grains was estimated. Thisfrequency was computed by counting the number of high aspect ratiograins which appear over a fixed lineal distance. At least about 30 betagrains/mm had a high aspect ratio. According to one toughening theory,the main toughening mechanism in a hipped silicon nitride-ceramic iscrack deflection, and crack deflection is a function of the frequency of"rods" within the ceramic. Since the aspect frequency estimates aresomewhat analogous to rod frequency, toughness was plotted as a functionof aspect frequency. See FIG. 4. Regression analysis of thetoughness-aspect frequency correlation resulted in a correlationcoefficient of about 0.97, indicating a strong correlation betweentoughness and aspect frequency.

The ceramics of the present invention may be used in conventionalsilicon nitride applications, including high temperature applicationssuch as the AGT rotor.

We claim:
 1. A process for producing a high tensile strength siliconnitride ceramic, comprising the steps of:a) milling a batch of siliconnitride powder in water to produce a slurry, b) passing the slurrythrough a filter having a pore size of between about 2 and about 10microns, c) concentrating the slurry to a solids content of betweenabout 70 and about 72 w/o, d) passing the slurry through a magneticseparator, e) casting the slurry at a CRS rate of less than about 10 mm²/min to produce a green body comprising at least about 80 w/o siliconnitride, f) drying the green body to less than about 3 w/o water, g)degassing the green body, h) sintering the green body to yield a denseceramic comprising at least about 80 w/o silicon nitride, i) machiningthe dense ceramic, and j) heat treating the dense ceramic to yield ahigh tensile strength silicon nitride ceramic.